Hot rolled high strength steel strip

Abstract

The present invention is based on the following components by weight: 0.02-0.13% C, 1.20-3.50% Mn, 0.10-1.00% Si, 0.05-1.0% Mo, 0.02-0.70% Al, 0.04-0.25% Ti, 0.010% or less N, 0.02% or less P, 0.01% or less S, 0.005% or less B, and, if necessary, (1.5% or less Cu, 0.70% or less Cr, 0.50% or less Ni, 0.30% or less V, 0.10% or less Nb) % ferrite and at most 5% martensite by volume, the steel strip having a misorientation distribution (MOD) index of at least 0.65 at quarter thickness, a maximum of 55% area fraction having a kernel mean misorientation (KAM) of 0-1, at least 45% area fraction having a kernel mean misorientation (KAM) of 1-5, and a thickness of at least 1400 mm. -1 The present invention relates to a hot rolled high strength steel strip having a total grain boundary length (ΣGB[5°-65°]) of at least 950 MPa, said steel strip having at least the following mechanical properties: an ultimate tensile strength (Rm) of at least 950 MPa, a total elongation (A50) of at least 10% and a hole expansion ratio (λ) value of at least 40%. The present invention also relates to said hot rolled steel strip and to an automotive part comprising said hot rolled steel strip.

Description

[Technical field]

[0001] The present invention relates to hot rolled high strength steel strip, in particular hot rolled high strength steel strip suitable for use in automotive parts. The present invention also relates to a method for producing said hot rolled high strength steel strip. Furthermore, the present invention relates to automotive parts comprising the hot rolled high strength steel strip. [Background technology]

[0002] It is well known that as the strength of hot rolled (HR) steel strip increases, the workability decreases. The main area of ​​application of HR steel in transportation and automotive applications is chassis and suspension (C&S). Other areas include frame rails of trucks, bumper beams or battery boxes of electric vehicles. The typical thickness of HR steel used for these applications is less than 4.5 mm. Thicker gauge HR steel strip, such as 12 mm or less, is used in engineering applications such as crane booms or transportation applications for the frames of large trucks. From the perspective of weight reduction, it is imperative that higher strength steels should be used in the above applications in order to reduce the gauge of the steel strip. Therefore, ultra-high strength steels, usually with an ultimate tensile strength (Rm) of more than 950 MPa, would be useful for this purpose. These applications of HR steel require mechanical properties that are difficult to match. In addition to high strength, the steel strip should also have excellent workability, for example to make parts by cold working, since this is an energy-efficient manufacturing route compared to hot working. Additionally, good impact and fracture toughness or energy absorption capacity is also required for applications such as bumper beams, battery housings, crane booms or frame rails. For assembly of parts, good weldability is also required. However, as the tensile strength of steel increases, the formability parameter decreases. Formability is a general term for steel sheets that is viewed as a combination of material behavior during several mechanical operations such as stretching, bending, drawing and flanging. Depending on the part geometry, any one or a combination of two or more attributes of the material are important during sheet metal fabrication.

[0003] To reduce the weight of parts, a common approach is to apply high strength steel and reduce the thickness of the steel strip used to reduce the weight. However, this can result in a loss of stiffness, which is undesirable in some applications of automotive parts. The inherent loss of stiffness due to the reduction in the thickness of the steel strip used to manufacture automotive parts can be regained by optimizing the part geometry, for example by making deeper flanges or flanges with increased degrees of stretching or bending. To enable the increase in part stiffness through geometry optimization, the high strength steel strip requires good processability in terms of tensile elongation and hole expansion ability.

[0004] Single-phase precipitation strengthened ferritic high strength steels that break away from the traditional constraint between global workability (e.g. tensile elongation) and local workability (e.g. hole expansion ability or HEC) with high levels of both workability modes can have low fracture toughness values ​​and increased edge crack sensitivity under certain conditions, e.g. under compression, tending to unstable brittle fracture behavior in shear compromising shear edge fatigue and increased sensitivity to delamination.

[0005] Patent document EP1616970A1 describes a process of reheating a steel slab containing, by weight, 0.04-0.15% C, 1.5% or less Si, 0.5-1.6% Mn, 0.04% or less P, 0.005% or less S, 0.04% or less Al, 0.03-0.15% Ti, 0.03-0.5% Mo (all by mass), and the balance being Fe and unavoidable impurities, at a temperature in the range of 1150-1300°C; A method for producing a high strength hot rolled steel sheet is disclosed, which comprises the steps of hot rolling a heated steel slab at a finishing temperature equal to or higher than the Ar3 transformation temperature into a hot rolled steel sheet; first cooling the hot rolled steel sheet to a temperature in the range of 700-850°C with an average cooling rate of 20°C / s or higher; holding the first cooled steel sheet at a temperature of 680°C or higher for more than 1 second; second cooling the steel sheet to a temperature of 550°C or lower with an average cooling rate of 30°C / s or higher, followed by coiling the steel sheet. Preferably, the hot rolled steel sheet is first cooled to a temperature in the range of 700-850°C and to a temperature in the range of (SRT / 3+300)-(SRT / 8+700)°C, where SRT represents the reheat temperature of the steel slab. In the examples of EP1616970A1, this indeed leads to a first hold at a temperature in the range of 667-860°C. The processing conditions are such that a microstructure is obtained consisting of ferrite containing precipitates, a second phase of bainite and / or martensite, and other phases, the percentage of ferrite containing precipitates being 40-95% and the percentage of other phases being 5% or less.

[0006] Patent document EP1338665A1 discloses a method for producing a high-strength hot-rolled steel sheet, the method including the steps of: producing a steel slab consisting essentially of, by weight, 0.06% or less C, 0.5% or less Si, 0.5-2.0% Mn, 0.06% or less P, 0.005% or less S, 0.1% or less Al, 0.006% or less N, 0.05-0.6% Mo, 0.02-0.10% Ti, and the balance Fe, and satisfying the formula 0.8≦(C / 12) / [(Ti / 48)+(Mo / 96)]≦1.3; hot rolling the steel slab at a temperature equal to or higher than the Ar3 transformation point to produce a hot-rolled steel sheet; and coiling the hot-rolled steel sheet at a temperature of 550-700°C. The treatment conditions are such that a microstructure consisting essentially of a matrix of a single ferritic phase and fine precipitates which are composite carbides containing Ti and Mo is obtained, the fine precipitates having a grain size of less than 10 nm and a grain size of 5×10 4 / m 3 The particles are dispersed in the matrix in the above number per unit volume.

[0007] There is a demand for hot rolled high strength steel strip that has high formability and low crack sensitivity. Summary of the Invention

[0008] As recognized herein, with respect to any description of an alloy composition or preferred alloy composition, all percentages are by weight unless otherwise specified.

[0009] As used herein, the term "about" to describe a compositional range or amount of an alloying addition means that the actual amount of the alloying addition may vary from the nominal intended amount due to factors such as standard processing variations, as one of ordinary skill in the art would understand.

[0010] As used herein, the terms "up to" and "up to about" expressly include, but are not limited to, the possibility of 0% by weight of the particular alloying element it represents. For example, 0.03% Cr or less may include the composition of steel strip having no Cr at all.

[0011] It is an object of the present invention to provide a hot rolled high strength steel strip having high total elongation (A50 or A80) together with high hole-expansion capacity.

[0012] It is an object of the present invention to provide a hot rolled high strength steel strip having high total elongation (A50 or A80) together with high hole expansion capacity and low crack sensitivity, in particular low edge-crack sensitivity.

[0013] It is another object of the present invention to provide a method for producing hot rolled high strength steel strip having an improved balance of total elongation, hole expansion ability and low crack sensitivity, especially low edge crack sensitivity.

[0014] These and other objects and further advantages are achieved by the following in weight percent: 0.02 to 0.13 wt% C; 1.20 to 3.50 wt.% Mn; 0.10 to 1.0 weight percent Si; 0.05 to 1.0 wt% Mo; 0.02 to 0.70 wt% Al; 0.04 to 0.25 weight percent Ti; 0.010% by weight or less N (100 ppm), preferably 0.0065% by weight or less N (65 ppm); 0.02 wt. % or less P, preferably 0.015 wt. % or less P; 0.01% by weight or less of S, preferably 0.0025% by weight or less of S (25 ppm); 0.0050% by weight or less of B (50 ppm), preferably 0.0030% by weight or less of B (30 ppm); as needed, % Cu or less, preferably 0.6% Cu or less, more preferably 0.10% Cu or less, 0.70 wt.% or less Cr, preferably 0.40 wt.% or less Cr, more preferably 0.25 wt.% or less Cr; 0.50 wt.% or less Ni, preferably 0.3 wt.% or less Ni, more preferably 0.10 wt.% or less Ni, 0.30 wt.% or less of V, preferably 0.20 wt.% or less of V, and 0.10% by weight or less of Nb, preferably 0.03% by weight or less of Nb One or more elements selected from the group consisting of The balance is Fe and unavoidable impurities resulting from the iron and steelmaking process. A hot rolled high strength steel strip having a composition consisting of: The steel strip has a microstructure of at least 95% by volume ferrite and at most 5% by volume martensite, preferably at least 0.2% by volume martensite; Steel strips are a Misorientation Distribution (MOD) index of at least 0.65 at quarter thickness; A maximum of 55% area fraction with Kernel Average Misorientation (KAM) between 0 and 1; at least 45% area fraction with a kernel mean misorientation (KAM) between 1 and 5; A total grain boundary length (ΣGB[5°~65°]) of at least 1400 1 / mm having The steel strip shall have at least the following mechanical properties: an ultimate tensile strength (Rm) of at least 950 MPa, preferably at least 960 MPa, more preferably at least 980 MPa; A total elongation (A50) of at least 10%, preferably at least 14%, and A hole expansion ratio (λ) value of at least 40% This is accomplished by the present invention which provides a hot rolled high strength steel strip having a

[0015] According to the present invention, it has been found that hot rolled steel strip having these narrow alloy composition ranges together with the microstructure provides an improved balance of high strength (Rm), total elongation (A50) and hole expansion ratio. It is also an important discovery of the present invention that the microstructure of the steel strip reduces edge crack sensitivity in processing operations. Reduced crack sensitivity, especially reduced edge crack sensitivity, can be objectively expressed in terms of average total crack length (ATCL) according to the measurement method described herein. The steel strip according to the present invention has an ATCL of less than 45 mm in preferred embodiments and in the best case less than 40 mm. [Brief description of the drawings]

[0016] [Figure 1] 1A-1E show some features of a method for determining the average total crack length (ATCL) of a steel strip product using the general cylindrical deep drawing test for drawn cups. [Diagram 2] FIG. 2A shows a schematic diagram of a hot rolling mill for processing thick cast steel slabs, and FIG. 2B shows a thin slab casting facility with a direct rolling mill. DETAILED DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0017] In one embodiment of the steel strip the yield strength (Rp) is at least 800 MPa, preferably at least 850 MPa. In one embodiment of the steel strip the steel strip has a yield strength (Rp) of at most 960 MPa, preferably at most 950 MPa.

[0018] The steel strip has an ultimate tensile strength (Rm) of at least 950 MPa. In one embodiment of the steel strip, the steel strip has an ultimate tensile strength of at least 960 MPa, preferably at least 980 MPa.

[0019] The steel strip has a total elongation (A50) of at least 10%, preferably at least 14%.

[0020] The required microstructure of the steel strip of the present invention is achieved by a narrow compositional range and by careful control of the manufacturing process, the accelerated cooling of the steel strip at the run-out table (ROT) and especially the narrow operating window for the coiling temperature (CT).

[0021] The microstructure of the steel strip according to the invention consists of at least 95% by volume of ferrite, precipitation strengthened by carbide precipitates of titanium and molybdenum, and optionally vanadium and / or niobium, and at most 5% by volume of martensite, the balance being unavoidable inclusions, totalling 100% by volume. In one embodiment, the microstructure has at least 98% by volume of ferrite, more preferably at least 99% by volume. According to the invention, the presence of some martensite is considered to be favorable in order to blunt any crack tip and reduce crack propagation. Indeed, the martensite may contain some traces of retained austenite, but retained austenite is preferably absent. The microstructure of the steel strip preferably has at least 0.2% by volume, more preferably at least 0.3% by volume of martensite. In a preferred embodiment, the steel strip has at most 3% by volume, more preferably at most 2% by volume of martensite.

[0022] The overall microstructural texture of the steel strip is further characterized by a sufficiently high Misorientation Distribution (MOD) index of at least 0.65, preferably at least 0.80, more preferably at least 0.83 and in the best case at least 0.90, the MOD index being measured at a quarter thickness of the steel strip.

[0023] The overall microstructural texture of the steel strip is further characterized by an area fraction of at most 55%, more preferably at most 52%, having a Kernel Mean Misorientation (KAM) between 0 and 1, and an area fraction of at least 45%, more preferably at least 48%, having a Kernel Mean Misorientation (KAM) between 1 and 5.

[0024] The overall microstructure of the steel strip is mm / mm 2 The steel is further characterized by a (5°-65°) grain boundary length (ΣGB) of at least 1400 per mm, i.e., at least 1400 1 / mm, preferably at least 1500 1 / mm. A higher ΣGB indicates a higher degree of grain refinement resulting in a desired lower crack sensitivity, as well as suppression or arrest of crack propagation, e.g., under compression, resulting in a desired shorter ATCL.

[0025] Carbon is present in an amount of 0.02-0.13 wt%. To achieve sufficient strength, a suitable minimum C content is 0.04 wt%, and in a preferred embodiment, at least 0.070 wt%. In a preferred embodiment, the C content is at most 0.12 wt%, which is beneficial for suppressing the effect of cooling rate dependency on the homogeneity of the final microstructure and for promoting high hole expansion capacity. Furthermore, C is an essential element for achieving precipitation strengthening in combination with carbide-forming microalloy elements such as titanium, niobium (if added) or vanadium (if added) and for scavenging C to suppress cementite formation in the final microstructure. By optimizing the other alloying elements, including Ti, Nb and / or V, it is possible to obtain a nearly homogeneous ferritic or bainite-ferritic microstructure that is substantially free of cementite.

[0026] To achieve sufficient hardenability and grain refinement, the steel strip has Mn in the range of 1.20 wt.% to 3.50 wt.%. In one embodiment, for an improved balance of strength, corrosion resistance, fracture toughness and edge crack sensitivity, the Mn content is in the range of 1.40 wt.% to 2.40 wt.%. To obtain sufficient grain refinement for improved fracture toughness and reduced crack sensitivity, the Mn content is preferably at least 1.50 wt.%. In one embodiment, the Mn content is at most 2.20 wt.%, more preferably at most 2.0 wt.%. Too high a Mn content may result in segregation during casting, which adversely affects the desired balance of properties.

[0027] Silicon is present in an amount of 0.10-1.0 wt% to improve the strength of the steel by substitutional solid solution strengthening of the iron lattice. Furthermore, Si is beneficial to suppress carbide precipitation (cementite and other carbides). However, when higher amounts of Si are used, the weldability and coatability of the steel deteriorates, so the amount of Si is preferably at most 0.95 wt%, in a preferred embodiment at most 0.70 wt%, more preferably at most 0.60 wt%.

[0028] Aluminum behaves in a manner equivalent to Si in the steel according to the invention. Aluminum slows down the carbide precipitation kinetics and inhibits the formation of cementite. If Al is less than 0.02 wt.%, the effect of inhibiting carbide formation is negligible. Values ​​of aluminum less than 0.02 wt.% are believed to be residuals from the deoxidation step during steelmaking, and therefore a minimum value of about 0.02 wt.% is preferred. On the other hand, if Al is more than 0.70 wt.%, excessive oxides may be formed during thermomechanical processing of the steel (slab reheating, hot rolling, coiling, etc.). Furthermore, Al increases the ferrite to austenite transformation temperature, necessitating the need for hot rolling of the steel at higher temperatures to finish hot rolling in the austenite phase, since intercritical ferrite appears at lower temperatures. Higher amounts of oxidation may occur at higher temperatures. These oxide scales are detrimental to hot rolling, pickling, coating and the overall surface appearance. Furthermore, the rolling force during hot rolling increases to such a level that, in combination with the presence of Si, Al above 0.70 wt% makes the steel very brittle and more difficult to hot roll. Therefore, Al in the present invention is present in an amount ranging from 0.02 to 0.70 wt%, preferably 0.02 to 0.60 wt%, more preferably 0.030 to 0.50 wt%.

[0029] Titanium is another essential alloying element, and is present in an amount of 0.04-0.25 wt.% to provide hardenability and to act as a carbide former that inhibits the formation of cementite while providing precipitation strengthening through the formation of small Ti-based carbides. However, Ti also combines with N, S, and C to form nitrides and carbon sulfides, depending on the specific chemical composition of the steel. Thus, at least 0.04 wt.% Ti is present to combine with substantially all the N and S in the steel, and to have sufficient excess Ti to combine with the C in the steel. If more than 0.25 wt.% Ti is present, coarse Ti nitrides, carbonitrides, and carbides may form that are difficult to dissolve during reheating of the slab before hot rolling. Furthermore, these coarse Ti nitrides, carbonitrides, and carbides result in a decrease in the hole expansion ability of the steel. Preferably, 0.07-0.22 wt.% Ti is present. In one embodiment, the Ti content does not exceed 0.20 wt.%, preferably does not exceed 0.15 wt.%, more preferably does not exceed 0.13 wt.%.

[0030] The steel strip has Mo as an intended alloying element present in the range of about 0.05% to 1.0% by weight for hardenability. In one embodiment, the Mo content is at least about 0.10% by weight, preferably at least about 0.20% by weight, more preferably at least 0.25% by weight, and most preferably at least 0.30% by weight. In one embodiment, the Mo content does not exceed about 0.80% by weight, preferably does not exceed about 0.55% by weight, and more preferably does not exceed 0.40% by weight. The presence of Mo may also improve the weldability of the steel strip.

[0031] In order to improve the thermal stability of the carbide precipitates and thus reduce the loss of precipitation strengthening due to precipitate coarsening resulting from thermal aging (e.g. thermal aging during coil cooling of the steel or during any subsequent thermal heat treatment (e.g. during coating or welding)), the atomic ratio of Mo to the total amount of the trace alloy elements Nb, Ti and V, expressed in weight % according to the following formula, is preferably maintained in the range of 0.75 to 1.25.

[0032]

number

[0033] Nitrogen, sulfur and phosphorus are residual elements present in steel as a result of the steel making and refining process. Their amounts are limited to 0.01 wt% or less S, 0.02 wt% or less P and 0.010 wt% or less N. Amounts greater than these are detrimental to mechanical properties, workability, toughness and weldability. In one embodiment, P is present at 0.015 wt% or less. In one embodiment, S is present at 0.005 wt% or less, more preferably 0.0025 wt% or less. N forms titanium nitrides with Ti and acts as a dispersoid for austenite grain size control during reheating. However, too much N results in excessively coarse TiN particles that may impair hole expansion ability. Preferably, the N content is 0.0065 wt% (65 ppm) or less. A suitable minimum N content is 0.0010 wt% (10 ppm).

[0034] Boron is not required to obtain the desired balance of properties of the steel strip, but may be present at up to 0.005% by weight, thus up to 50 ppm, preferably up to 0.0030% by weight (30 ppm). B is very effective for enhancing the hardenability of the steel, which means that no or only little pro-eutectic ferrite is formed, allowing the use of lower carbon contents and / or lower cooling rates in the run-out table. B is also a very suitable alloying element for increasing the yield strength.

[0035] Cr may be added to the steel strip in an amount of about 0.70% by weight, preferably up to 0.25% by weight, to improve the hardenability of the steel. In a preferred embodiment, the steel strip has Cr as an acceptable impurity element. In practice, this means that Cr may be present up to 0.10%, preferably up to 0.050%, more preferably up to 0.03%, as this may impair the corrosion resistance of the uncoated steel strip substrate.

[0036] Copper, when present at 1.5 wt.% or less, improves the strength of the steel by both solid solution strengthening as well as precipitation hardening from copper precipitates. In one embodiment, the Cu content does not exceed 0.6 wt.%. In one embodiment, Cu is not added as a purposeful alloying element and may be present at 0.10 wt.% or less, more preferably 0.05 wt.% or less. Nickel, at about 0.50 wt.% or less, preferably 0.30 wt.% or less, improves impact toughness and combats any hot shortness that may occur during hot working of the steel strip due to the presence of copper. In one embodiment, Ni is not added as a purposeful alloying element and may be present at 0.3 wt.% or less, preferably 0.10 wt.% or less, more preferably 0.050 wt.% or less.

[0037] Vanadium may be present in the steel in an amount of about 0.30 wt% or less, preferably about 0.20 wt% or less, more preferably 0.17 wt% or less. However, V is a relatively expensive alloying element that is primarily used to replace Ti for its precipitation strengthening effect and to avoid the formation of cementite by the production of vanadium carbides. In a preferred embodiment, V is purposefully added in an amount of at least 0.05 wt%, more preferably at least 0.08 wt%. A combined addition of Ti and V is preferred because Ti provides a catalytic effect on V precipitation, making it more effective.

[0038] Niobium may be present in the steel at up to 0.10 wt.%. Nb improves the strength of steel, partly by precipitation hardening and mainly by grain refinement. However, at high amounts of Nb, these effects saturate. Therefore, preferably at most 0.08 wt.%, more preferably at most 0.060 wt.% Nb is present. Since Nb is a rather expensive alloying element, in one embodiment, Nb is not purposefully added but is present as an impurity element, not exceeding 0.03 wt.%, preferably not exceeding 0.02 wt.%, more preferably not exceeding 0.0050 wt.%. In the most preferred embodiment, Nb is an unavoidable impurity resulting from the iron and steel making process. It has been found that by maintaining Ti at a sufficiently high level, the use of Nb can be avoided while still achieving the desired balance of workability, mechanical properties, and improved fracture or edge crack properties. Furthermore, Nb has a high tendency to lead to segregation due to centre-line segregation and the formation of coarse NbC particles, therefore avoiding the use of Nb leads to improved hole expansion capacity and improved shear edge quality.

[0039] Furthermore, it has been found that the formation of cementite can be suppressed and the preferred formation of a small fraction of martensite + retained austenite in the microstructure can be better controlled if the amounts of the carbide forming elements Ti, Nb, V and Mo, expressed in weight percent, satisfy the following formula:

[0040]

number

[0041] In the above formula, Ti_sol is defined as the amount of free Ti in solution,

number

[0042] Preferably, the lower limit of the above formula is 0.75, preferably 0.80. In one embodiment, to further suppress the presence of cementite and / or to control the amount of martensite, the upper limit is preferably 1.8, more preferably 1.5. In a preferred embodiment, the above formula is in the range of 0.9 to 1.1.

[0043] In one embodiment the steel strip comprises 0.02-0.13 wt.% C, 1.20-2.0 wt.% Mn, 0.10-0.60 wt.% Si, 0.01-0.70 wt.% Al, 0.04-0.25 wt.% Ti, 0.05-0.80 wt.% Mo, 0.10% or less Cr, preferably 0.050% or less Cr, 0.010% or less N, 0.02% or less P, preferably 0.015% or less P, and 0.01% or less. The composition is S, preferably 0.0025% by weight or less S, 0.0050% by weight or less B, and optionally one or more elements selected from the group consisting of (0.10% by weight or less Cu, 0.10% by weight or less Ni, 0.30% by weight or less V, 0.10% by weight or less Nb, preferably 0.03% by weight or less Nb), with the balance being Fe and unavoidable impurities resulting from the iron and steel making process, more preferred ranges being described in the present specification and claims.

[0044] The 0.2% offset proof strength or yield strength (Rp), ultimate tensile strength (Rm), uniform elongation (Ag) and tensile elongation (A50) were measured at room temperature using A50 specimen geometry at a quasi-static strain rate of 3×10 -4 seconds -1 ) tensile test parallel to the rolling direction according to EN 10002-1 / 150 6892-1. The geometry of the tensile specimen was 50 mm gauge length in the rolling direction, 12.5 mm width and thickness according to the final gauge. The strength of the steel at 0.2% offset strain is measured as the yield strength (Rp or YS). The ratio of the yield strength to the ultimate tensile strength (Rp / Rm) is expressed as the yield ratio.

[0045] The stretch-flangeability or hole expansion capacity (HEC) of steel strips was determined by hole expansion tests. Specimens with dimensions 90 mm × 90 mm × final thickness of strip were cut from the as-coiled steel. A hole with a diameter of 10 mm was punched in the centre of the specimen and hole expansion tests were carried out according to the ISO / TS16630:2003(E) standard. The hole expansion tests of the samples were carried out by upward burring. A 60° conical punch was pushed up from below and the hole diameter d was determined when a through thickness crack was formed. f was measured. The hole expansion ratio λ is d 0 = 10 mm and was calculated using the following formula:

[0046]

number

[0047] For all of the above mechanical tests, at least three specimens were tested for each condition, and the average values ​​are reported herein.

[0048] The average total crack length (ATCL) is used to evaluate the susceptibility of crack formation in a situation similar to industrial applications. The ATCL parameter is determined in a common laboratory cylindrical deep drawing test with a punch, a drawing die and a blank holder. In this test, the punch has a diameter of 50 mm and a punch radius of 7 mm. The die has a diameter of 62 mm and a radius of 8 mm. A schematic of this setup is shown in Figure 1B. The inner diameter is large enough to allow the edge of the formed cup to move freely. The clearance, i.e., the distance between the punch wall and the die wall, is 6 mm. The blank holder force is set to 50 kN. The blank is rectangular and 90 x 90 mm. The four corners of the rectangle are cut by 10 mm in the direction of the two diagonals, as shown in Figure 1A. During the initial stage of the cylindrical deep drawing test, the four regions at the edge of the blank undergo plastic deformation due to high local compressive stresses during drawing. This results in localized wrinkling of the edges. At the end of the test and upon release of the blank holder force, the four compressed regions are subjected to reverse loading due to spring back as these regions of the formed cylindrical cup begin to lose contact with the blank holder. This reverse loading due to spring back can result in the nucleation and growth of cracks in the four regions of the compressed and wrinkled drawn cup (see, for example, FIG. 1E). The cracks may penetrate the entire thickness of the steel strip and be visible on both sides of the drawn cup, i.e., inside and outside, or may be visible only on one of the two sides of the drawn cup. The length of all visible cracks on the inside and outside of the four compressed edges of the deep drawn cup (shown in FIG. 1D) is measured using a 10x magnifying glass with a scale grid. The sum of the lengths of all observable cracks on the inside and outside of the cup wall is averaged for three of the four drawn cups and reported as the average total crack length (ATCL), expressed in mm.

[0049] The microstructure of the steel strips was analyzed by Electron Back Scatter Diffraction (EBSD). EBSD is a well-known technique in the art, which in turn allows the quantification of the area or volume fraction of the various components. EBSD measurements were performed on cross sections parallel to the rolling direction (RD-ND plane), encapsulated in conductive resin and mechanically polished to 1 μm. A final polishing step was performed with colloidal silica (OPS) to obtain a completely distortion-free surface.

[0050] The scanning electron microscope (SEM) used for EBSD measurements was a Zeiss Ultra55 instrument equipped with a field emission gun (FEG-SEM) and an EDAX PEGASUS XM4 HIKARI EBSD system. EBSD scans were collected on the RD-ND plane of the sheet at 1 / 4 thickness. The sample was placed at an angle of 70° in the SEM. The accelerating voltage was 15 kV and the high current option was switched on. An aperture of 120 μm was used and the normal working distance during scanning was 17 mm. Dynamic focus correction was used during scanning to compensate for the high tilt angle of the sample.

[0051] EBSD scans were captured using TexSEM Laboratories (TSL) software "Orientation Imaging Microscopy (OIM) Data Collection version 7.2". Typically, the following data collection settings were used: Hikari camera at 5 × 5 binning combined with background subtraction (standard mode). In all cases, the scan area was located at ¼ sample thickness, and care was taken to avoid non-metallic inclusions in the scan area as much as possible. In all cases, the EBSD scan size was 100 × 100 μm, the step size was 0.1 μm, and the scan speed was approximately 100 frames / s. Fe(α) and Fe(γ) were used to index the Kikuchi patterns. The hough settings used during data collection were: binned pattern size approximately 96, theta set size 1, rho fraction approximately 90, maximum peak count 10, minimum peak count 5, hough type setting classic, hough resolution setting low, butterfly convolution mask 9 × 9, peak symmetry 0.5, minimum peak magnitude 10, maximum peak distance 20.

[0052] EBSD scans were evaluated using TSL OIM analysis software version 8.0x64[12-14-16]. Typically, the data set was rotated 90° with respect to the RD axis to obtain a scan with the proper orientation relative to the measurement direction. A standard grain dilation clean-up was performed (Grain Tolerance Angle (GTA): 5°, minimum grain size: 5 pixels, criteria that grains must contain multiple rows for a single dilation iteration clean-up). This was followed by a pseudo-symmetry clean-up (GTA5, axis angle 30°@111) applied.

[0053] EBSD Image Quality (IQ) maps were used to determine the amount of martensite. Areas with low IQ were identified as MS regions. For a given experimental condition, the low IQ threshold was typically about 0.4 of the peak maximum location in the IQ histogram. However, to prevent grain boundaries from granular bainite or upper bainitic areas in the martensite area fraction from being included in the martensite area fraction, the low IQ threshold was manually checked for each scan.

[0054] For the calculation of the EBSD Kernel Average Misorientation (KAM) map, the fifth nearest neighbor was used with a maximum misorientation of 5° (all points in the kernel were used for the KAM calculation). The Kernel Average Misorientation is considered a signature of the ferrite type since it is a measure of the internal dislocation density. Regions with relatively low internal dislocation density primarily coincide with regions with KAM values ​​between 0 and 1°. Regions with relatively high internal dislocation density primarily coincide with regions with KAM values ​​between 1 and 5°.

[0055] 100×100μm area (0.01mm) on the RD-ND plane at 1 / 4 thickness 2 ) from the EBSD scan, furthermore, from 5°~15°(ΣGB 5~15 ) and 15°~65°(ΣGB 15~65 The misorientation angle of the grain boundaries was measured along with the sum of all grain boundary lengths (ΣGB). The value of ΣGB is expressed in mm -1 It is expressed in units of ΣGB and is a measure of the density of high-angle grain boundaries. High-angle grain boundaries are effective in arresting crack propagation. Therefore, an increase in the value of ΣGB is beneficial for improving fracture toughness and reducing crack sensitivity.

[0056] One aspect of the present invention, as described herein and in the claims, relates to a method for producing steel strip, the method comprising the steps of: after casting the slab, reheating the solidified slab to a temperature between 1050°C and 1260°C, preferably for a time period of at least 30 minutes, more preferably at least 60 minutes, and hot rolling the slab or, after casting the slab or strip, hot rolling the slab or strip; Hot rolling the steel slab or strip and finishing the hot rolling at a finish rolling temperature between 820°C and 940°C, preferably between 850°C and 940°C, most preferably between 850°C and 920°C and higher than the Ar3 temperature of the steel (the finish rolling temperature (FRT) is higher than the Ar3 temperature of the steel, Ar3 being the temperature at which austenite starts to transform to ferrite during cooling. As known in the art, the Ar3 temperature can be determined by the following formula: Ar3 = 910℃-203×[C] 1 / 2 +44.7×[Si]-30×[Mn]+31.5×[Mo] can be calculated according to ); accelerated cooling of the hot rolled steel strip at a run-out table cooling rate of 20-250°C / s, preferably 40-200°C / s, to a run-out table temperature of 560°C-620°C; coiling the hot rolled and cooled strip at a temperature between 550°C and 600°C, preferably between 550°C and 595°C, more preferably between 550°C and 590°C; further cooling the coiled hot rolled steel strip to ambient temperature; · Process of pickling hot rolled steel strip; Optionally providing the hot rolled steel strip with a metallic coating layer, preferably selected from the group comprising Zn, Zn-based alloy and Al-based alloy layers, to improve corrosion resistance during use; in the order given. The metal coating layer is preferably formed by heat-to-coat or hot-dip coating.

[0057] The manufacturing methods described herein provide a desired microstructure that provides a targeted improved balance of workability, mechanical properties, and fracture properties. The invention is also embodied in steel strip produced by the methods described herein, having the above microstructure and improved balance of workability, mechanical properties, and fracture properties.

[0058] The invention is not limited by the casting method. The steel can be cast as conventional thick slabs with a cast thickness of 150 mm to 350 mm, typically 225 mm to 250 mm, and as thin slabs with a cast thickness of 50 mm to 150 mm in a direct strip plant. Schematic examples of the method including a conventional hot rolling mill and a thin slab caster / direct rolling mill are shown in Figures 2A and 2B, respectively. In the case of conventional thick slab casting, reheating of the slab is necessary to reheat the slab from ambient temperature (usually the thick cast slabs have been cooled from the casting temperature to ambient temperature in the slab yard) and to homogenize the slab in terms of composition, therefore the reheating temperature should be above about 1050°C in order to dissolve any precipitates if microalloy elements are present and to bring the slab to a temperature such that the final hot rolling in the finishing mill is still feasible even with FRT>Ar3. Often this requires a (slab) reheat temperature of 1050°C to about 1260°C. In the case of thin slab casting, the cast slab undergoes homogenization treatment in a homogenization furnace immediately after casting of the thin slab, the homogenization temperature should be above about 1050°C, typically about 1100-1160°C. This also suppresses the formation of any precipitates of micro-alloying elements (if any) present, and brings the thin slab to a temperature such that final hot rolling in the finishing mill is still feasible even with FRT>Ar3. According to the invention, the reheat or homogenization time for the thin slab casting route is preferably 30 minutes or more.

[0059] Hot rolling of steels must be performed in the austenite phase to control the final microstructure. At the industrial scale of rolling FRT, the FRT should be kept above the Ar3 temperature. In a preferred embodiment, the FRT is above (Ar3+30°C), e.g., typically above 850°C, to avoid hot rolling being performed locally at a temperature lower than Ar3 at the colder edge or tail of the stop. It should not be too high in the austenite region, since a lower FRT promotes more austenite distortion and thus contributes to increased grain refinement and increased ΣGB. An FRT above 950°C results in increased edge crack sensitivity. Furthermore, a not too high FRT can improve toughness (e.g., {332} <113> ) to promote the final microstructural texture that is beneficial to the grains and to prevent the grains from being adversely affected (e.g., {001} <110> For this reason, the FRT should not exceed about 940°C, preferably not exceed about 920°C, and more preferably not exceed about 910°C.

[0060] After hot rolling, the steel strip is accelerated cooled on a run-out table (ROT) to a temperature of 560°C-620°C. The accelerated cooling rate is desirable to suppress the recovery and loss of internal stored energy in austenite to promote grain refinement and increase in ΣGB of the final microstructure. The cooling rate should be high enough to avoid the austenite to ferrite phase transformation at high temperatures and to favorably promote the austenite to ferrite phase transformation at the relatively low temperatures of about 560-630°C on the run-out table. Increasing the cooling rate promotes grain refinement, increase in ΣGB, and thus increase fracture toughness and decrease crack sensitivity. Increasing the cooling rate also suppresses texture randomization, and thus improves toughness (e.g., {332} <113> ) to suppress the loss of strength of these textures developed from the deformed austenite. The cooling rates referred to herein are not critical run-out table cooling rates (ROT-CR) as long as they exceed through-thickness of the steel strip from a microstructural point of view. However, an unnecessarily high ROT-CR can affect the planarity of the strip after cooling and cause control problems for stopping at the correct cooling stop temperature, therefore a suitable maximum ROT-CR is about 250°C / s, preferably about 200°C / s, more preferably about 150°C / s. The practical ROT-CR range is about 20-100°C / s, more preferably about 40-100°C / s, which can be achieved by air cooling, laminar cooling or water jet cooling, depending on the thickness of the steel strip. For practical reasons, the run-out table cooling rate (ROT-CR) is defined as the average cooling rate on the surface of the steel strip.

[0061] The hot rolled steel strip is then coiled at a temperature of 550°C to 600°C, preferably 550°C to 600°C, more preferably 560°C to 600°C. In one embodiment the hot rolled strip is coiled at a temperature not exceeding 595°C, more preferably not exceeding 590°C. The coiling temperature of the steel strip is a critical process parameter for achieving the required microstructure in the steel strip that provides the improved balance of mechanical properties described herein.

[0062] If the coiling temperature is too low, the kinetics of the precipitates will be insufficient, resulting in low strength levels. If the coiling temperature is too high, the grain refinement will be insufficient, leading to reduced fracture toughness and increased edge crack sensitivity. Also, too high a coiling temperature will not promote martensite as a second phase constituent in the final microstructure. Also, too high a coiling temperature will reduce the MOD index, increase the fraction KAM0-1, and make the fraction KAM1-5 too small. During coil cooling, some further precipitation may occur, as well as some further phase transformations. Undesirably, the precipitates once formed may coarsen during coil cooling. The alloy composition in combination with the claimed coiling temperature suppresses this phenomenon. This coiling temperature helps to promote the refinement of the grains of the ferrite formed during coiling or coil cooling, suppress the coarsening of the precipitates that strengthen the ferrite matrix, and also promote the formation of a small amount of martensite.

[0063] In patent document EP1616970A1, holding temperatures below 680°C result in insufficient driving force for the ferrite transformation and subsequently in a too low proportion of ferrite with precipitates. In the present case, the austenite to ferrite phase transformation is enhanced at temperatures in the range below 680°C in order to achieve increased grain refinement and increased ΣGB for increased fracture toughness and reduced crack sensitivity while still having sufficient kinetics for precipitation. Furthermore, patent document EP1616970A1 discloses that after holding the steel strip at a temperature of 680°C or higher for more than 1 second, it is necessary to apply a second cooling with an average cooling rate of 30°C / s or higher, preferably 50°C / s or higher, to a coiling temperature of 550°C or lower, preferably 450°C or lower, more preferably 350°C or lower, and coiling to form a second phase of bainite and / or martensite and to suppress the formation of other phases to less than 5% by volume. In the present case, the coiling temperature is rather high, between 550°C and 600°C, preferably between 550°C and 595°C, more preferably between 550°C and 590°C, to allow the austenite to ferrite phase transformation to continue at a relatively low temperature and promote fine-grained ferrite, which is precipitate strengthened with carbide precipitates containing Ti and Mo and optionally Nb and / or V. The grain refinement and increase in ΣGB provide increased fracture toughness and reduced (edge) crack sensitivity. Coiling below 550°C results in insufficient ferrite formation and loss of precipitation. Furthermore, this may result in too high a martensite fraction.

[0064] Patent document EP1338665A1 discloses coiling temperatures in the range of 550-700°C and discloses that steels having a tensile strength of at least 950 MPa and a hole expansion capacity of at least 40% were all produced using coiling temperatures above 600°C. The finish rolling temperatures for all these steels were in the range of 880-930°C. According to the present invention it has been found that too high coiling temperatures lead, inter alia, to an increased crack sensitivity, in particular an increased edge crack sensitivity, as well as a decrease in fracture toughness.

[0065] After the steel strip has cooled to room temperature, the oxides (scale) on the hot rolled steel strip are removed either by pickling in an acid solution (e.g., HCl) at warm temperatures (80-120°C) or by a combination of pickling and mechanical brushing of the strip surface. This step is necessary to make the steel strip surface suitable for direct use as uncoated hot rolled steel or suitable for coating treatments, if required for corrosion resistance.

[0066] In one embodiment, the hot rolled steel strip has a thickness in the range of about 1.5-8 mm, more preferably about 1.8-6 mm, and most preferably about 1.8-4 mm.

[0067] The hot rolled steel strip product may be a bare or uncoated product or may be coated on one or both major faces of the hot rolled steel strip product, typically with a coating weight of about 100 g / m2 per side of the steel strip. 2 or less, preferably about 50 g / m per side 2 The following thin metal coating layers may be provided: The metal coating is preferably selected from the group comprising aluminium alloy coatings (e.g. Al-Si alloys or Al-Zn alloys), zinc coatings and zinc alloy coatings (e.g. Zn-Al alloys, Zn-Mg alloys, Zn-Fe alloys, Zn-Al-Mg alloys or Zn-Mg-Al alloys).

[0068] The composition of the zinc or zinc alloy coating layer is not limited. The coating layer can be formed in a variety of ways, but hot dip galvanizing is preferred, using a standard Gl coating bath. The Zn-based coating layer may include a Zn alloy containing Al as an alloying element. The preferred zinc bath composition includes about 0.10-0.35 wt. % Al, the balance being zinc and unavoidable impurities.

[0069] Other zinc coating layers may be applied. An example is a zinc alloy coating according to patent document WO2008 / 102009A1 (herein incorporated by reference), in particular a zinc alloy coating layer consisting of 0.3-4.0 wt.% Mg and 0.05-6.0 wt.% Al, preferably 0.1-5.0% Al, optionally up to 0.2 wt.% of one or more additional elements, unavoidable impurities and the balance zinc. A preferred Zn bath containing Mg and Al as main alloying elements has the composition: 0.5-3.8 wt.% Al, 0.5-3.0 wt.% Mg, optionally up to 0.2 wt.% of one or more additional elements; the balance zinc and unavoidable impurities. The additional elements, which are typically added in small amounts of less than 0.2 wt.%, may be selected from the group comprising Pb, Sb, Ti, Ca, Mn, Sn, La, Ce, Cr, Ni, Zr and Bi. Pb, Sn, Bi and Sb are usually added to form spangles. The total amount of the additional elements in the zinc alloy is preferably at most 0.2% by weight, more preferably at most 0.1% by weight. These small amounts of additional elements do not change the properties of the coating or bath to a significant extent for normal applications. When one or more additional elements are present in the coating, each additional element is preferably present in an amount of 0.02% by weight or less, and more preferably each additional element is present in an amount of 0.01% by weight or less. The additional elements are usually only added to prevent dross formation in baths containing molten zinc alloys for hot dip galvanizing or to form spangles in the coating layer.

[0070] In another embodiment, the metal coating comprises a (commercially pure) aluminum layer or an aluminum alloy layer. A typical metal bath for hot-dipping such an aluminum layer comprises aluminum alloyed with silicon, for example about 8-11% by weight of silicon, and up to 4% by weight of iron, optionally up to 0.2% by weight of one or more additional elements (e.g. calcium), unavoidable impurities, balance aluminum. Silicon is present to prevent the formation of thick iron-metallic intermetallic layers that reduce adhesion and workability. Iron is preferably present in an amount of 1-4% by weight, more preferably at least 2% by weight.

[0071] One aspect of the present invention relates to a galvanized steel strip obtainable by hot-dip galvanizing a hot-rolled high strength steel strip according to the present invention.

[0072] One aspect of the present invention relates to automotive parts, in particular automotive chassis parts, which comprise or are made from the hot rolled high strength steel strip according to the present invention and which benefit, inter alia, from an improved balance of strength, workability, increased fracture toughness and reduced (edge) crack sensitivity. The steel strip can be formed into automotive parts by cold, warm and hot working operations as known in the art. Automotive parts include suspension arms, reinforcements, body-in-white frame materials as side members, seat frames, seat rails, bumper beams, battery boxes for electric vehicles, all parts with complex shapes. By using the hot rolled high strength steel strip, these parts can be produced with high quality, cost-effectiveness and high yield. The high strength hot rolled steel products according to the present invention can also be used in engineering applications.

[0073] The invention will now be described with the aid of the following non-limiting drawings.

[0074] 1A-1E show some features of a method for determining the average total crack length (ATCL) of a steel strip product using the general cylindrical deep drawing test for drawn cups. FIG. 2A shows a schematic diagram of a hot rolling mill for processing thick cast steel slabs, and FIG. 2B shows a thin slab casting facility with a direct rolling mill. EXAMPLES

[0075] The present invention will now be described with reference to non-limiting comparative examples and inventive examples.

[0076] Five steel ingots of the inventive examples (Inv.) Steels A to E and the comparative example (Comp.) Steel F, measuring 320 x 100 x 100 mm, were cast by melting the charges in a vacuum induction furnace. The chemical composition of these steels is given in Table 1, which also gives the following two ratios A and B:

[0077]

number

[0078]

number

[0079] All ingots were reheated at 1240°C for 1 hour and rough rolled to a thickness of 35 mm. The strips were then reheated again to 1220°C for 40 minutes and hot rolled to a final thickness of about 3.2-3.6 mm in 5 rolling passes with FRT above 870°C and below 910°C, which is the austenite phase field, for all these steels. After the finishing rolling pass, the hot rolled steels were rolled according to the start temperature run-out table (T START The steel is then moved to a run-out table (T ) in the range of 840-880°C, and then heated with a mixture of water and air from the austenite phase field to the final temperature of the ferrite phase field on a run-out table (T ) in the range of 565-615°C. ENDThe steel was then actively cooled at 300°C / s at rates ranging from 30 to 70°C / s. The steel was then transferred to a furnace for repeated slow coil cooling at furnace temperatures (CT - coiling temperature) of 540°C, 580°C and 610°C (see Table 2).

[0080] The hot rolled sheets were sandblasted to remove the oxide layer before testing for mechanical properties, hole expansion capacity and average total crack length according to the methods described herein. The microstructure of the hot rolled strip was determined using EBSG according to the methods described herein. The results of these tests are shown in Table 2 as a function of the applied processing parameters.

[0081] [Table 1]

[0082] From the results in Table 2, it can be seen that Steel 17, which has a composition outside the claimed range and which has been processed according to the invention (CT 580°C), has too low an Rm. Although increasing the coiling temperature (CT 610°C) increases the Rm, the edge crack sensitivity of Steel 16 is significantly reduced. This reduction is most likely the result of a microstructure with too low an MOD index and an insufficient degree of grain refinement reflected by ΣGB. In the case of Steel 18, a coiling temperature that is too low (CT 540°C) leads to an even lower Rm compared to Steel 17.

[0083] Steels 1-3 (Alloy A) all have a composition according to the invention, and Steel 2 was processed (CT 580°C) according to the invention. Steel 2 offers a very good balance of high strength (Rm 1011 MPa), high elongation (A50 15.8%), good hole expansion ratio (HEC 42%) and very good edge crack resistance (ATCL 28.9mm). This is due to the composition, processing and microstructure. The microstructure is characterized by a very high degree of grain refinement (ΣGB 1854 1 / mm) and a high MOD index (1.20). In the case of Steel 1, an increase in CT to 610°C still gives a high Rm, but the edge crack sensitivity decreases to unacceptable levels (ATCL 49.3mm). This is a result of microstructural changes reflected, inter alia, by a too low MOD index (0.52), too high KAM0-1 and too low KAM1-5, and by a too small grain refinement reflected by a too low ΣGB. For Steel 3, the reduction in CT to 540°C leads to a significant decrease in Rp and Rm.

[0084] A similar trend to that of Steels 1 to 3 (Alloy A) is observed in Steels 4 to 6 (Alloy B).

[0085] Steels 7-9 (Alloy C) all have compositions according to the invention with purposely added V. Steel 8 processed according to the invention at CT 580°C offers an excellent balance of high strength (Rm 1005MPa), high elongation (A50 14.7%), excellent hole expansion ratio and very good edge crack resistance (ATCL 34.8mm). This is due to the composition, processing and microstructure. The microstructure is characterized by a very high degree of grain refinement. The microstructure has about 2.6 vol.% martensite, which is believed to contribute to the excellent edge crack resistance represented by the ATCL of 34.8mm. In the case of Steel 7, increasing the CT to 610°C slightly increases Rm and Rp, but the edge crack sensitivity of Steel 7 decreases significantly. This decrease is likely the result of the changes in the microstructure reflected by the MOD index and KAM values, as well as an insufficient degree of grain refinement reflected by ΣGB (ΣGB 1134 1 / mm). In the case of Steel 9, a coiling temperature that is too low (CT 540° C.) results in a low Rm and a significantly lower Rp as well as a hole expansion ratio that is too low compared to Steel 8.

[0086] Alloy D (Steels 10-12) have increased Al content compared to Alloy B (Steels 4-6). Comparison of Steels 5 and 11 (both processed according to the present invention) shows that increased strength (Rm) is obtained with the addition of Al. Steel 11 also exhibits favorable microstructure as indicated by the MOD index, KAM value, and a sufficient degree of grain refinement as indicated by ΣGB, thereby providing a desired balance of properties.

[0087] Alloy E (steels 13-15) have a purposeful addition of V compared to alloy D (steels 10-12). Steel 14 offers a further reduced edge crack sensitivity compared to steel 11. Comparison of Steel 14 with Steel 15 shows that a lower coiling temperature (CT 540°C) results in, among other things, a significantly reduced Rp. Comparison of Steel 13 with Steel 14 also shows that in the case of Steel 13, an increase in CT to 610°C still gives a high Rm, but the edge crack sensitivity is reduced to an unacceptable level (ATCL 56.3 mm). This is reflected, among other things, by a too low MOD index, too high KAM0-1 and low KAM1-5, as well as a low ΣGB.

[0088] While the present invention has been described to the greatest extent possible, it will be apparent to those skilled in the art that numerous changes and modifications can be made without departing from the spirit or scope of the invention as set forth herein.

[0089] [Table 2]

Claims

1. In weight percent, 0.02 to 0.13 wt. % C; 1.20 to 3.50 wt. % Mn; 0.10 to 1.0 wt. % Si; 0.05 to 1.0 wt. % Mo; 0.02 to 0.70 wt. % Al; 0.04 to 0.25 wt. % Ti; 0.010 wt.% or less of N; 0.02 wt. % or less of P; 0.01 wt.% or less of S; 0.005 wt. % or less of B; as needed, 1.5 wt.% or less of Cu, 0.70 wt.% or less of Cr, 0.50 wt.% or less of Ni, 0.30 wt. % or less of V, and 0.10 wt% or less Nb one or more elements selected from the group consisting of: The balance is Fe and unavoidable impurities.

1. A hot rolled high strength steel strip having a composition consisting of: the steel strip has a microstructure of at least 95% by volume ferrite and at most 5% by volume martensite; The steel strip a misorientation distribution (MOD) index of at least 0.65 at quarter thickness; an area fraction of at most 55% with a kernel average misorientation (KAM) between 0 and 1; an area fraction of at least 45% having a kernel mean misorientation (KAM) of 1 to 5; A total grain boundary length (ΣGB[5°-65°]) of at least 1400 1 / mm and The steel strip has at least the following mechanical properties: an ultimate tensile strength (Rm) of at least 950 MPa; a total elongation (A50) of at least 10%, and Hole expansion ratio (λ) value of at least 40% The hot rolled high strength steel strip having

2. 2. The hot rolled high strength steel strip of claim 1, wherein the steel strip has at least 98% by volume ferrite.

3. The hot rolled high strength steel strip of claim 2, wherein the steel strip has at least 99% by volume of ferrite.

4. 2. The hot rolled high strength steel strip of claim 1, wherein the overall microstructure of the steel strip has 0.2 to 4 volume percent martensite.

5. A hot rolled high strength steel strip as claimed in claim 4, wherein the overall microstructure of the steel strip has 0.2 to 3 volume percent martensite.

6. A hot rolled high strength steel strip as claimed in claim 5, wherein the overall microstructure of the steel strip has 0.2 to 2 volume percent martensite.

7. 2. The hot rolled high strength steel strip of claim 1, wherein the microstructure has a misorientation distribution (MOD) index of at least 0.

80.

8. The hot rolled high strength steel strip of claim 7, wherein the microstructure has a misorientation distribution (MOD) index of at least 0.

83.

9. 2. The hot rolled high strength steel strip of claim 1, wherein the microstructure has a total grain boundary length (ΣGB[5°-65°]) of at least 1500 1 / mm.

10. The composition is in the range of 0.45 to 2.2, [Equation 1] [Wherein Ti_sol is [Equation 2] is defined as follows.] 2. The hot rolled high strength steel strip of claim 1, having

11. The composition is in the range of 0.55 to 2.1, [Equation 3] [In the formula, Ti_sol has the same meaning as defined above.] 11. The hot rolled high strength steel strip of claim 10, having

12. 2. The hot rolled high strength steel strip of claim 1, wherein the steel strip has a Mo content in the range of 0.10 to 0.80 wt.%.

13. The hot rolled high strength steel strip of claim 12, wherein the steel strip has a Mo content in the range of 0.25 to 0.80 wt. %.

14. Hot rolled high strength steel strip according to claim 13, wherein the steel strip has a Mo content in the range of 0.30 to 0.80 wt.%.

15. 2. The hot rolled high strength steel strip of claim 1, wherein the steel strip has a Cr content of 0.10 wt. % or less.

16. Hot rolled high strength steel strip according to claim 15, wherein the steel strip has a Cr content of 0.050 wt.% or less.

17. Hot rolled high strength steel strip according to claim 16, wherein the steel strip has a Cr content of 0.03 wt. % or less.

18. 2. The hot rolled high strength steel strip of claim 1, wherein the steel strip has a V content in the range of 0.05 to 0.30 wt. %.

19. Hot rolled high strength steel strip according to claim 18, wherein the steel strip has a V content in the range of 0.08 to 0.20 wt. %.

20. 2. The hot rolled high strength steel strip of claim 1, wherein the steel strip has a Mn content in the range of 1.40 to 2.40 wt.%.

21. Hot rolled high strength steel strip according to claim 20, wherein the steel strip has a Mn content in the range of 1.40 to 2.20 wt.%.

22. 2. The hot rolled high strength steel strip of claim 1, wherein the steel strip has an average total crack length (ATCL) of less than 45 mm.

23. The hot rolled high strength steel strip of claim 22, wherein the steel strip has an average total crack length (ATCL) of less than 40 mm.

24. The hot rolled high strength steel strip of claim 1 , wherein the steel strip comprises a metallic coating layer.

25. The hot rolled high strength steel strip of claim 24, wherein the metal coating layer is selected from the group comprising a Zn layer, a Zn-based alloy layer and an Al-based alloy layer.

26. 10. A method for producing hot rolled high strength steel strip according to claim 1, comprising the steps of: - after casting the slab, reheating the solidified slab to a temperature of 1050°C to 1260°C and hot rolling the slab, or hot rolling the slab or strip immediately after casting the slab or strip; Hot rolling the steel slab or strip and finishing the hot rolling at a finishing rolling temperature of 820°C to 940°C and above the Ar3 temperature of the steel; Accelerated cooling of the hot rolled steel strip at a run-out cooling rate of 20-250°C / s to a run-out table temperature of 560°C-620°C; Coiling the hot-rolled and cooled strip at a temperature of 550°C to 600°C; Further cooling the coiled hot rolled steel strip to ambient temperature; - Process of pickling hot rolled steel strip The method comprising:

27. The method of claim 26, wherein the finish rolling temperature is between 850°C and 940°C and higher than the Ar3 temperature of the steel.

28. The method of claim 26, wherein the runout cooling rate is 40 to 200°C / sec.

29. The method of claim 26, wherein the temperature in the step of coiling the hot-rolled and cooled strip is between 550°C and 595°C.

30. The method of claim 29, wherein the temperature in the process of coiling the hot-rolled and cooled strip is between 550°C and 590°C.

31. 27. The method of claim 26, wherein the pickled hot rolled steel strip comprises a metallic coating layer.

32. The method of claim 31, wherein the metal coating layer is selected from the group consisting of a Zn layer, a Zn-based alloy layer, and an Al-based alloy layer.

33. The method described in claim 31, wherein the metal coating layer is a metal coating layer formed by heat-to-coat or hot-dip plating.

34. 34. An automotive part comprising a hot rolled high strength steel strip according to any one of claims 1 to 25 or obtainable by a method according to any one of claims 26 to 33.

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